CN113195750B

CN113195750B - High-strength steel material and method for producing same

Info

Publication number: CN113195750B
Application number: CN201980082754.7A
Authority: CN
Inventors: 朱尼·塔斯特; 铁波·皮卡莱宁; 汤米·利马泰宁; 卡蒂·里廷基
Original assignee: SSAB Technology AB
Current assignee: SSAB Technology AB
Priority date: 2018-12-11
Filing date: 2019-12-11
Publication date: 2023-05-23
Anticipated expiration: 2039-12-11
Also published as: KR20210099627A; CN113195750A; PL3666911T3; EP3666911B1; US20220106654A1; US11505841B2; DK3666911T3; JP2022512191A; EP3666911A1; WO2020120563A1; ES2895456T3

Abstract

A high strength steel comprising a composition consisting of, in weight percent: 0.02% to 0.05% C, 0.1% to 0.6% Si, 1.1% to 2.0% Mn, 0.01% to 0.15% Al, 0.01% to 0.08% Nb, 0.5% or less Cu, 0.5% or less Cr, 0.7% or less Ni, 0.03% or less Ti, 0.1% or less Mo, 0.1% or less V, 0.0005% or less B, 0.015% or less P, 0.005% or less S, the balance being Fe and unavoidable impurities, wherein the steel has a microstructure comprising a matrix consisting of, in volume percent: 40% to 80% of quasi-polygonal ferrite, 20% to 40% of polygonal ferrite, 20% or less of bainite, and the balance of pearlite and martensite of 20% or less. The steel has a yield strength of at least 400MPa, an ultimate tensile strength of at least 500MPa, and a tensile strength of at least 34J/cm at a temperature in the range of-50 ℃ to-100 DEG C ² Charpy-V impact toughness.

Description

High-strength steel material and method for producing same

Technical Field

The invention relates to a high-strength ultralow-carbon steel material which can be used for manufacturing pressure vessels, gas conveying pipelines and building materials. The invention also relates to a method for manufacturing the high-strength ultralow-carbon steel.

Background

The general trend in steel development is toward a combination of higher strength and low temperature impact toughness, and good weldability. Conventional and standard thick plate pressure vessel steels (e.g., ASTM a537CL 2) are traditionally manufactured using carbon contents of 0.1 to 0.2 weight percent (wt.%) to obtain adequate strength. Due to the high carbon content, the weldability of these steels decreases, the toughness is poor and the resistance to Hydrogen Induced Cracking (HIC) is low. Therefore, for good formability, low Carbon Equivalent (CE), low impact transition temperature, good Crack Tip Opening Displacement (CTOD), and high resistance to Post Weld Heat Treatment (PWHT), it is necessary to reduce the carbon content in the steel.

Low carbon (C) steels have been developed in which carbon is not the primary strength source, as high carbon concentrations may lead to poor weldability and weld toughness. Furthermore, high carbon concentrations may impair the impact toughness of the steel. One of the first studies on very low carbon steel was McEvily et al from ford motor company in 1967. They showed that 0.04C-3.0Ni-3.0Mo-0.05Nb gave a yield strength of about 700MPa and a transition temperature of about-75 ℃. However, this composition is highly alloyed, seeking more economical alloying elements that provide equivalent performance.

To compensate for the strength loss due to low carbon content, alloy design concepts have been based on advanced use of cost-effective micro-alloying elements such as niobium (Nb), titanium (Ti), vanadium (V) and boron (B) in combination with moderate amounts of other alloying elements such as manganese (Mn), silicon (Si), chromium (Cr), molybdenum (Mo) and copper (Cu) to improve the hardenability of austenite. The mature use of the aforementioned combination of (micro) alloying elements in combination with low carbon content can produce steels with yield strengths ranging from 500MPa to 900 MPa. These (micro) alloying elements increase strength by microstructural refinement, precipitation hardening and solid solution strengthening, and strengthening by microstructural modification.

Typically, low carbon microalloyed steel is processed by a thermo-mechanical control process (TMCP), which traditionally consists of three stages. In the first rough rolling stage, the austenite grain size is refined due to repeated cycles of the recrystallization process. In the second controlled rolling stage, the austenite is deformed in the non-recrystallization temperature range, which brings about significant refinement of the final ferrite microstructure. In the final stage, accelerated cooling may be employed to further refine the resulting ferrite grain size while inhibiting the formation of polygonal ferrite and facilitating the formation of low temperature transformation products (such as different types of bainite). Therefore, these low carbon microalloyed steels with high strength are commonly referred to as Low Carbon Bainite (LCB) steels. The combination of low carbon and ultra fine ferrite grain size provides a good combination of strength and toughness, and has good weldability due to low carbon and low alloy content.

The combination of TMCP and the use of (micro) alloys has an impact on the development of microstructure related to mechanical properties. In continuously cooled low carbon microalloyed steels, the predominant austenitic decomposition product is ferrite. However, a portion of the parent austenite may not be transformed, but may remain or partially transform at room temperature to form martensite-austenite (MA) micro-components. At very high cooling rates, even very low carbon steels with sufficient hardenability may transform into martensite.

The microstructure of LCB steels is often complex, consisting of a mixture of different ferrite morphologies, ranging from polygonal ferrite to lath martensite. Classification systems and terms proposed by the japanese steel association bainitic committee (ISIJ) can be used to characterize all possible ferrite morphologies formed in low carbon steels. A short description of all six ferrite morphologies is as follows.

1. Polygonal Ferrite (PF) exhibits substantially equiaxed grains and smooth boundaries.

2. Quasi-polygonal ferrite (QF) appears to have undulating boundaries of grains that may cross previous austenitic boundaries that contain dislocation substructures and occasional MA micro-components. This is also called bulk ferrite.

3. Weishi (Wei)

Ferrite (WF) exhibits elongated ferrite crystals with minimal dislocation substructure.

4. Granular Bainite (GB) exhibits a succession of elongated ferrite crystals (grains or equiaxed shapes), has low orientation obstruction and high dislocation density, and contains islands of MA components that are substantially equiaxed.

5. Bainitic Ferrite (BF), also known as Acicular Ferrite (AF), has a number of parallel ferrite laths or sheets separated by low angle boundaries and having a very high dislocation density. The MA component remaining between ferrite crystals has a needle-like morphology.

6. The dislocated cubic martensite exhibits a highly dislocated lath-like morphology, preserving the prior austenite boundaries.

EP 2484792 A1 relates to a low carbon steel having a three-phase microstructure consisting of 5% to 70% of bainite and 3% to 20% of MA components in terms of area fraction, the remainder being quasi-polygonal ferrite. In order to secure strength, the area fraction of the quasi-polygonal ferrite is preferably 10% or more. The 5% to 70% bainite ensures toughness of the substrate. The MA composition of 3% to 20% ensures a lower yield ratio and toughness of the substrate. The triphasic microstructure precludes the presence of polygonal ferrite or other microstructures. Low carbon steel has a low yield ratio, high strength, high toughness and excellent strain aging resistance. Low carbon steel is produced by a process comprising the steps of: heating to a temperature in the range 1000 ℃ to 1300 ℃; hot rolling at a final rolling temperature not lower than an Ar 3 transformation temperature, wherein a cumulative rolling reduction in an austenite non-recrystallization temperature range is 50% or more; accelerated cooling to a stop temperature of 500 ℃ to 680 ℃; and re-heated to a temperature of 550 to 750 ℃.

EP 2380997 A1 describes a low carbon steel for welded structures, which has excellent high temperature strength, low temperature toughness and suppressed weld crack parameters. The high temperature strength is ensured by co-addition of Cr and Nb, which contributes to transformation strengthening and precipitation strengthening. A low carbon steel comprising a bainitic structure is manufactured by a method comprising the steps of: heating to a temperature in the range 1000 ℃ to 1300 ℃, preferably 1050 ℃ to 1250 ℃; hot rolling at a final hot rolling temperature of 800 ℃ or more, preferably 800 ℃ or more; and accelerated cooling to a stop temperature of 550 ℃ or less, preferably 520 ℃ to 300 ℃.

JP 2007119861 (a) or JP 2007277679 (a) also relates to a low carbon steel for welded structures, which has excellent high temperature strength, low temperature toughness and suppressed weld crack parameters. Low carbon steel comprising a martensite-austenite mixed phase (i.e. MA component) is manufactured by a method comprising the steps of: heating to a temperature in the range 1000 ℃ to 1300 ℃; performing hot rolling at a final rolling temperature of 750 ℃ or more, wherein the cumulative rolling reduction in the austenite non-recrystallization temperature range is 30% or more; and accelerated cooling to a stop temperature of 350 ℃ or less. It is noted in the specification that when the accelerated cooling is stopped at a temperature of 230 ℃, the hardness difference between the surface and the center of the steel sheet having a thickness of 50mm becomes extremely large, so that the bendability and the hole expansibility may be adversely affected.

KR 20030054424 (a) relates to a non-heat treated low carbon steel having high weldability, high toughness and high tensile strength of more than 600 MPa. It was found that it is necessary to prevent polygonal ferrite from forming in austenite grain boundaries to ensure strength. In order to obtain excellent toughness, it is necessary to control the cumulative rolling compression in the range of 30% to 60% in the austenite non-recrystallization temperature region. If the cumulative rolling reduction in the austenite non-recrystallization temperature range is less than 30%, the low temperature toughness cannot be effectively improved. If the cumulative rolling compression in the austenite non-recrystallization temperature range excessively increases and exceeds 60%, the effect of lowering the transformation temperature is saturated, and the anisotropy increases, so that a plate deformation problem occurs during use.

The object of the present invention is to further develop a high strength low carbon steel and a method for manufacturing the same, such that a new steel material having mechanical properties without damage and economic advantages can be obtained.

Disclosure of Invention

In view of the prior art, an object of the present invention is to solve the problem of providing a high strength low carbon steel having excellent low temperature impact toughness, bendability/formability and weldability, which are required in applications such as fusion welding of pressure vessels and structures. This problem can be solved by combining a cost-effective (micro) alloy design with a cost-effective TMCP procedure, which combination yields a metallographic microstructure mainly comprising quasi-polygonal ferrite.

In a first aspect, the present invention provides a high strength steel comprising a composition or composition consisting of, in weight percent (wt.%):

c0.02-0.05, preferably 0.03-0.045

Si 0.1 to 0.6, preferably 0.2 to 0.6, more preferably 0.3 to 0.5

Mn 1.1-2.0, preferably 1.35-1.8

Al 0.01-0.15, preferably 0.02-0.06

Nb is 0.01-0.08, preferably 0.025-0.05

Cu is less than or equal to 0.5, preferably 0.15 to 0.35

Cr is less than or equal to 0.5, preferably 0.1 to 0.25

Ni is less than or equal to 0.7, preferably 0.1 to 0.25

Ti is less than or equal to 0.03, preferably 0.005 to 0.03

Mo≤0.1

V is not more than 0.1, preferably not more than 0.05

B≤0.0005

P is not more than 0.015, preferably not more than 0.012

S≤0.005

The balance being iron and unavoidable impurities.

The steel is low alloyed and cost effective alloying elements such as C, si, mn, al and Nb. Other elements (e.g., cu, cr, ni, ti, mo, V and B) may be present as residual levels that are not intentionally added. The difference between the residue content and the unavoidable impurities is that the residue content is a controlled amount of alloying elements that are not regarded as impurities. The residual content normally controlled by the industrial process has no essential effect on the alloy.

Preferably, the steel comprises nonmetallic inclusions having an average inclusion size in the range of 1 μm to 4 μm in diameter, and wherein 95% of the inclusions have a diameter of less than 4 μm.

In a second aspect, the present invention provides a method for manufacturing a high strength steel material, the method comprising the steps of:

-heating a steel billet having the composition or composition according to claim 1 to a temperature in the range of 950 ℃ to 1350 ℃;

-hot rolling the heated steel slab in a plurality of hot rolling passes, wherein

i. Subjecting the steel blank to a first plurality of rolling passes at a temperature above the austenite non-recrystallization temperature;

cooling the steel blank from step (i) to a temperature below the austenite non-recrystallization temperature,

subjecting the steel blank from step (ii) to a second plurality of controlled rolling passes at a temperature below the austenite non-recrystallization temperature, wherein the controlled rolling passes have a reduction ratio of at least 1.5, preferably 2.0, more preferably 2.5, and wherein the final rolling temperature is in the range 800 ℃ to 880 ℃;

-accelerating the continuous cooling to a temperature below 230 ℃ at a cooling rate of at least 5 ℃/s.

At a temperature below the austenite non-recrystallization temperature T _nr The controlled rolling pass subjected to at the temperature of (a) leads to the accumulation of austenitic deformation which leads to the formation of elongated grains and deformed strips. Grain boundaries and bands of deformation can act as nucleation sites for austenite to ferrite (gamma-alpha) transformation. As austenite grains elongate, the grain boundaries also become closer and closer, increasing nucleation density. In combination with the high nucleation rate resulting from accelerated continuous cooling, the process ultimately results in ultra-fine ferrite grain sizes.

After accelerating the continuous cooling, an additional tempering step is optionally performed at a temperature in the range of 580 ℃ to 650 ℃ for 0.5 to 1 hour. The additional tempering step may optionally be induction tempering at a temperature typically in the range 580 ℃ to 700 ℃ for 1 to 60 minutes.

Preferably, the cumulative compression ratio of the hot rolling is in the range of 4.0 to 35.

The processing parameters must be strictly controlled to improve the mechanical properties, in particular toughness, the main parameters involved being the heating temperature, the cumulative compression ratio of the controlled rolling passes at a temperature lower than the austenite non-recrystallization temperature, the final rolling temperature and the accelerated continuous cooling stop temperature.

The steel is a strip or plate having a thickness of 6mm to 65mm, preferably 10mm to 45 mm.

The steel product obtained has a microstructure comprising a matrix consisting of, in volume percent (vol.%):

quasi-polygonal ferrite 40-80

Polygonal ferrite and bainite 20-60

Pearlite and martensite are 20 or less, preferably 5 or less, more preferably 2 or less.

Preferably, the microstructure includes polygonal ferrite in an amount of 20 to 40% by volume.

Preferably, the microstructure comprises bainite in an amount of 20 vol% or less.

A good combination of strength and toughness is associated with a microstructure based on quasi-polygonal ferrite. The steel has the following mechanical properties:

a yield strength of at least 400Mpa, preferably at least 415Mpa, more preferably in the range of 415Mpa to 650Mpa;

an ultimate tensile strength of at least 500MPa, preferably in the range 500MPa to 690MPa, more preferably in the range 550MPa to 690 MPa;

at a temperature in the range of-50 ℃ to-100 ℃, a Charpy-V impact toughness of at least 34J/cm ² Preferably at least 150J/cm ² More preferably at least 300J/cm ² 。

The steel material exhibits excellent bendability or formability. The steel has a minimum bending radius in the longitudinal or transverse direction of 5.0t or less, preferably 3.0t or less, more preferably 0.5t, and wherein t is the thickness of the steel strip or sheet.

Thus, improvements in properties such as low temperature impact toughness, bendability/formability, and weldability, HIC resistance, and PWHT resistance can be achieved. Post-weld heat treatment at a temperature in the range of 500 ℃ to 680 ℃ for 1 hour to 8 hours, or at a temperature in the range of 600 ℃ to 640 ℃ for 4 hours to 8 hours has little or no negative effect on the steel.

Drawings

Fig. 1 is a graph showing the Yield Strength (YS) of a production lot of 2000 tons of sheet.

Fig. 2 is a graph showing the Ultimate Tensile Strength (UTS) of a production lot of 2000 ton board.

Fig. 3 is a graph showing the Total Elongation (TEL) of a production lot of 2000 ton plates.

Fig. 4 is a graph showing impact toughness values at-45 ℃ (KV) for a production lot of 2000 ton board.

Fig. 5 is a graph showing the charpy-V impact toughness of panels having different thicknesses.

Fig. 6 is a graph showing the NACE TM 0284HIC test results for plates with different thicknesses.

Fig. 7 is a graph showing the mechanical properties (YS, UTS, TEL) of plates with different thickness under transport or PWHT conditions.

Fig. 8 is a graph showing the results of full thickness tensile tests for plates having thicknesses of 12mm, 25mm and 41 mm.

Fig. 9 is a graph showing impact toughness levels for plates having different thicknesses.

FIG. 10 is a graph showing the effect of rolling parameters on longitudinal Charpy-V impact toughness in a plate having a thickness of 25 mm.

FIG. 11 is a graph showing the effect of rolling parameters on longitudinal Charpy-V impact toughness in a plate having a thickness of 41 mm.

Fig. 12 illustrates the microstructure of the test specimen.

Detailed Description

The term "steel" is defined as an iron alloy comprising carbon (C).

The term "(micro) alloying element" is used to denote

Microalloying elements (MAE), such as niobium (Nb), titanium (Ti), vanadium (V) and boron (B); and/or

Medium content of alloying elements such as manganese (Mn), silicon (Si), chromium (Cr), molybdenum (Mo) and copper (Cu).

The term "nonmetallic inclusion" refers to the products of pollution, chemical reactions, and physical effects that occur during manufacturing. Nonmetallic inclusions include oxides, sulfides, nitrides, silicates, and phosphides.

The term "austenite non-recrystallization temperature" (T) _nr ) Defined as a temperature below which no complete static recrystallization of austenite occurs between rolling passesTemperature.

The term "Controlled Rolling (CR)" means that the rolling temperature (CR) is lower than the austenite non-recrystallization temperature (T _nr ) Is a hot rolling at a temperature of (a).

The term "compression ratio" refers to the ratio of thickness compression obtained by a rolling process. The compression ratio is calculated by dividing the thickness before the rolling process by the thickness after the rolling process. A compression ratio of 2.5 corresponds to a 60% reduction in thickness.

The term "controlled rolling ratio" means by being below T _nr Is subjected to controlled rolling at a temperature of (a) to obtain a compression ratio.

The term "cumulative compression ratio" means by being higher and lower than T _nr The total compression ratio obtained by hot rolling at the temperature of (a).

The term "Accelerated Continuous Cooling (ACC)" refers to a process of continuously accelerating cooling to a certain temperature at a certain cooling rate.

The term "Intermittent Accelerated Cooling (IAC)" refers to a process of performing accelerated cooling at a cooling rate in a certain temperature range and then air-cooling it to a temperature lower than the temperature range.

The term "ductile-brittle transition temperature (DBTT)" is defined as the lowest temperature at which steel has absorbed a specified amount of energy without cracking. At temperatures above the DBTT, steel can bend or deform like plastic at impact; whereas at temperatures below the DBTT, steels have a greater propensity to fracture or shatter upon impact.

The term "ultimate tensile strength (UTS, rm)" refers to the limit value at which the steel breaks under tension and is therefore the maximum tensile stress.

The term "yield strength (YS, rp _0.2 ) "refers to an offset yield strength of 0.2%, which is defined as the amount of stress that will result in a plastic strain of 0.2%.

The term "Total Elongation (TEL)" refers to the percentage of a material that can be stretched before breaking; the rough formability index is typically expressed as a percentage over a fixed gauge length of the measured extensometer. Two common gauge lengths are 50mm (a ₅₀ ) And 80mm (A) ₈₀ )。

The term "minimum bend radius (Ri)" is used to refer to the minimum bend radius that can be applied to a test sheet without cracking.

The term "bendability" refers to the ratio of Ri to sheet thickness (t).

The symbol "KV" refers to the absorption energy required to break a V-notch test piece of a prescribed shape and size when tested with a pendulum impact tester.

The alloy content of the steel and the processing parameters determine the microstructure, which in turn determines the mechanical properties of the steel.

Alloy design is one of the primary issues to be considered in developing steels with targeted mechanical properties. In general, it can be said that the lower the carbon content and the higher the target strength level, a higher content of the alternative (micro) alloying element is required in order to obtain an equivalent strength level.

The chemical composition is described in more detail below, wherein% of each component refers to weight percent.

The carbon C is used in a range of 0.02% to 0.05%.

Carbon alloying increases the strength of steel by solid solution strengthening, so the carbon content determines the strength level. Carbon contents less than 0.02% may result in insufficient strength. However, carbon has a detrimental effect on the weldability, weld toughness and impact toughness of the steel. Carbon also increases the DBTT. Therefore, the carbon content is set to not more than 0.05%.

Preferably, the carbon is used in the range of 0.03% to 0.045%.

The silicon Si is used in a range of 0.1% to 0.6%.

Silicon is an effective deoxidizer or sedative (k-ing agent) that removes oxygen from the melt during the steelmaking process. Silicon alloying increases strength by solid solution strengthening and hardness by increasing austenite hardenability. The presence of silicon may also stabilize the retained austenite. However, a silicon content higher than 0.6% may unnecessarily increase the Carbon Equivalent (CE) value, thereby impairing solderability. Furthermore, if silicon is excessively present, the surface quality may deteriorate.

Preferably, silicon is used in a range of 0.2% to 0.6%, and more preferably 0.3% to 0.5%.

The manganese Mn is used in a range of 1.1% to 2.0%.

Manganese is an essential element for improving the balance between strength and low temperature toughness. There appears to be a rough relationship between higher manganese content and higher strength levels. Manganese alloying increases strength by solid solution strengthening and hardness by increasing austenite hardenability. However, alloying with a manganese content exceeding 2.0% may unnecessarily increase CE value, thereby impairing solderability. If the manganese content is excessively high, hardenability of the steel increases, so that not only Heat Affected Zone (HAZ) toughness is deteriorated, but also center line segregation of the steel sheet is promoted, and as a result, low temperature toughness of the center of the steel sheet is impaired.

Preferably, manganese is used in a range of 1.35% to 1.8%.

The aluminum Al is used in a range of 0.01% to 0.15%.

Aluminum is an effective deoxidizer or sedative that removes oxygen from the melt during the steelmaking process. Aluminum may also remove nitrogen by forming stable AlN particles and provide refined grains that promote high toughness, especially at low temperatures. The aluminum also stabilizes the retained austenite. However, excessive aluminum may increase nonmetallic inclusions, thereby deteriorating cleanliness.

Preferably, the aluminum is used in a range of 0.02% to 0.06%.

The niobium Nb is used in the range of 0.01% to 0.08%.

Niobium forms carbides NbC and carbonitrides Nb (C, N). Niobium is considered to be the main grain refining element. Niobium promotes steel strengthening and toughening by four means:

i. by introducing fine Nb (C, N) precipitates, the austenite grain structure is refined due to the pinning effect of Nb (C, N) during the reheating and soaking stages at high temperature;

at high temperatures (> 1000 ℃), the recrystallization kinetics are retarded due to the niobium solute drag effect and the occurrence of recrystallization due to strain-induced precipitation is prevented at lower temperatures, thus contributing to the refinement of the microstructure;

precipitation strengthening during and/or after (or subsequent to) the gamma-alpha conversion; delay the phase change to a lower temperature, resulting in phase change hardening and toughening.

Niobium is a preferred alloying element in these steels because it promotes the formation of a quasi-polygonal ferrite/bainite grain microstructure rather than polygonal ferrite. However, since further increasing the content of niobium does not have a significant effect on further improving the strength and toughness, the addition amount of niobium should be limited to 0.08%. Niobium may be detrimental to HAZ toughness because niobium may promote the formation of coarse-grain upper bainite structures by forming relatively unstable TiNbN or TiNb (C, N) precipitates.

Preferably, niobium is used in the range of 0.025% to 0.05%.

The copper Cu is used in a range of 0.5% or less.

Copper may promote lower bainite structure, resulting in solid solution strengthening and contributing to precipitation strengthening. Copper also has the beneficial effect of combating HIC and Sulfide Stress Corrosion Cracking (SSCC). When added in excess, copper may deteriorate in-situ solderability and HAZ toughness. Therefore, the upper limit thereof is set to 0.5%.

Preferably, copper is used in a range of 0.15% to 0.35%.

The chromium Cr is used in a range of 0.5% or less.

Chromium, as a medium strength carbide forming element, increases the strength of the parent steel and weld joint, with a slight sacrifice in impact toughness. Chromium alloying increases strength and hardness by increasing austenite hardenability. However, if the chromium is used in an amount exceeding 0.5%, the HAZ toughness and the field weldability may be adversely affected.

Preferably, chromium is used in the range of 0.1% to 0.25%.

Nickel Ni was used in a range of 07% or less.

Nickel is an alloying element that improves austenite hardenability to increase strength without loss of toughness and/or HAZ toughness. However, if the nickel content exceeds 0.7%, the alloy cost increases too much without significant technical improvement. Excessive nickel may produce high viscosity iron oxide scale that deteriorates the surface quality of the steel. Higher nickel content can also negatively impact solderability due to increased CE values and crack susceptibility coefficients.

Preferably, nickel is used in the range of 0.1% to 0.25%.

The titanium Ti is used in a range of 0.03% or less.

By forming stable TiN together with NbC, adding titanium to incorporate free nitrogen detrimental to toughness, the growth of austenite grains in the reheating stage at high temperature can be effectively prevented. The TiN precipitates may further prevent coarsening of the crystal grains of the HAZ during welding, thereby improving toughness. Formation of TiN suppresses Fe ₂₃ C ₆ Thereby stimulating nucleation of polygonal ferrite. The formation of TiN also inhibits precipitation of BN, thereby allowing B to participate freely in its contribution to hardenability. For this purpose, the ratio Ti/N is at least 3.4. However, if the Ti content is too high, coarsening of TiN and precipitation hardening due to TiC develop, and low-temperature toughness may be deteriorated. Therefore, titanium must be limited to less than 0.03%, preferably less than 0.02%.

Preferably, ti is used in a range of 0.005% to 0.03%.

The molybdenum Mo is used in an amount of 0.1% or less.

Molybdenum has the effect of promoting a lower bainite structure while suppressing the formation of polygonal ferrite. Molybdenum alloying improves low temperature toughness and tempering resistance. The presence of molybdenum also increases strength and hardness by increasing austenite hardenability. In the case of boron alloying, molybdenum is typically required to ensure the effectiveness of the boron. Molybdenum, however, is not an economically acceptable alloying element. If the molybdenum is used in an amount of more than 0.1%, toughness may be deteriorated, thereby increasing the risk of brittleness. Excessive amounts of molybdenum may also reduce the effect of boron.

The vanadium V is used in an amount of 0.1% or less.

Vanadium has substantially the same but smaller effect than Nb. Vanadium is a strong carbide and nitride forming element, but V (C, N) can also be formed and its solubility in austenite is higher than Nb or Ti. Thus, vanadium alloying has the potential for dispersion and precipitation strengthening, as a large amount of vanadium dissolves and can precipitate in the ferrite. However, since polygonal ferrite is formed instead of bainite, the addition of vanadium exceeding 0.1% adversely affects weldability and hardenability.

Preferably, vanadium is used in an amount of 0.05% or less.

The boron B is used in an amount of 0.0005% or less.

Boron is a mature microalloying element for inhibiting the formation of diffusion transformation products such as polygonal ferrite, thereby promoting the formation of lower bainite structure. Effective boron alloying would require the presence of Ti to prevent BN formation. In the presence of boron, the Ti content can be reduced to less than 0.02%, which is very beneficial for low temperature toughness. However, when the boron content exceeds 0.0005%, the low-temperature toughness and the HAZ toughness deteriorate rapidly.

The unavoidable impurities may be phosphorus in an amount of 0.015% or less, preferably 0.012% or less; the sulfur content is 0.005% or less. Other unavoidable impurities may be nitrogen, hydrogen, oxygen, rare Earth Metals (REM), etc. In order to ensure excellent mechanical properties such as impact toughness, their content is limited.

Clean steelmaking processes are employed to minimize unavoidable impurities that may be present as nonmetallic inclusions. Nonmetallic inclusions disrupt the structural uniformity and thus their impact on mechanical and other properties can be significant. Nonmetallic inclusions may cause cracking and fatigue failure of the steel in deformation due to flattening, forging, and/or stamping. Therefore, the average inclusion size is generally limited to 1 μm to 4 μm, with 95% of the inclusions having a diameter of less than 4 μm.

The high strength steel may be a strip or plate of typical thickness of 6mm to 65mm, preferably 10mm to 45 mm.

The parameters of TMCP are adjusted to achieve the optimal microstructure with this chemical composition.

During the heating phase, the slab is heated to a discharge temperature in the range of 950 ℃ to 1350 ℃ (typically 1140 ℃), which is important for controlling the growth of austenite grains. The increase in heating temperature causes dissolution and coarsening of the microalloy precipitates, resulting in abnormal grain growth.

In the hot rolling stage, the slab is hot rolled, typically using 16-18 hot rolling passes, depending on the thickness of the slab and the final product. Preferably, at the end of the hot rolling stage, the cumulative compression ratio is in the range 4.0 to 35.

The first hot rolling process is carried out at a temperature above the austenite non-recrystallization temperature, and then at a temperature below T _nr Before controlled rolling at a temperature below T _nr Is set in the temperature range of (a).

Controlled rolling at a temperature below the austenite non-recrystallization temperature results in elongation of the austenite grains and creates starting sites for ferrite grains. Pancake-shaped austenite grains are formed, so that strain (i.e., dislocation) is accumulated in the austenite grains, and ferrite grain refinement can be promoted by acting as nucleation sites for austenite-to-ferrite transformation. A controlled rolling ratio of at least 1.5, preferably 2.0, more preferably 2.5 ensures a sufficient deformation of the austenite grains. A controlled rolling reduction of 2.5 can be achieved with 4 to 10 rolling passes, with a reduction of about 10.25% per pass. The most significant result of deformation in the austenite non-recrystallized region is an improvement in toughness performance. Surprisingly, the inventors found that increasing the controlled rolling compression ratio from 1.8 to 2.5 or higher can significantly reduce the transition temperature, thereby improving the low temperature impact toughness.

The final rolling temperature is typically in the range 800 ℃ to 880 ℃, which aids in the refinement of the microstructure.

Accelerated cooling of the hot rolled product to low at a cooling rate of at least 5 ℃/sAt a temperature of 230 ℃, preferably room temperature. At a value higher than Ar ₃ During rapid accelerated cooling from the temperature of (2) to the cooling stop temperature, refinement of ferrite grains is promoted. A low temperature transformation microstructure such as bainite is also formed in the accelerated cooling step.

Optionally, a subsequent step of heat treatment such as tempering or annealing is performed to fine tune the microstructure. Preferably, tempering is performed at a temperature in the range of 580 ℃ to 650 ℃ for 0.5 hours to 1 hour. An additional step of tempering may optionally be induction tempering for 1 to 60 minutes, typically at a temperature in the range 580 to 700 ℃.

During the accelerated continuous cooling, polygonal ferrite transformation occurs first at a reduced temperature, and then quasi-polygonal ferrite transformation, bainite transformation, and martensite transformation are sequentially performed. The final steel has a mixed microstructure based on quasi-polygonal ferrite. The microstructure comprises, in volume percent, 40% to 80% of quasi-polygonal ferrite and 20% to 60% of polygonal ferrite and bainite, the remainder being less than 20%, preferably less than 5%, more preferably 2% pearlite and martensite. Optionally, the microstructure comprises 20% to 40% polygonal ferrite by volume percent. Optionally, the microstructure comprises 20% or less by volume of bainite. Sometimes islands of MA component can be detected in the microstructure.

Good toughness and especially low DBTT of steel are often associated with high-density high-angle boundaries, which are often present in microstructures and are therefore beneficial because these boundaries can hinder crack propagation. The dominant microstructure of the quasi-polygonal ferrite is advantageous for forming high angle boundaries between the interfaces of the quasi-polygonal ferrite and the granular bainitic ferrite, while the formation of the quasi-polygonal ferrite eliminates the prior austenite grain boundaries in the microstructure.

The quasi-polygonal ferrite-based microstructure also reduces the size and proportion of MA micro-components, which are believed to be favorable nucleation sites for brittle fracture. The distribution of the MA component is limited to the granular bainitic ferrite portion of the microstructure.

If cracking microcracks are initiated near the MA micro-component, the propagation of the microcracks tends to passivate and temporarily stop due to the adjacent high angle boundaries. In order for the microcracks to reach a critical length beyond which they will propagate in an unstable manner, more energy is required to connect and link adjacent microcracks by, for example, short microcracks rotating in shear mode. Thus, steels with a quasi-polygonal ferrite-based microstructure have improved impact toughness and especially low DBTT.

The steel has a yield strength of at least 400MPa, preferably at least 415MPa, more preferably from 415MPa to 650MPa; the ultimate tensile strength is at least 500MPa, preferably in the range of 500MPa to 690MPa, more preferably in the range of 550MPa to 690 MPa. The steel has a Charpy-V impact toughness of at least 34J/cm at a temperature in the range of-50 ℃ to-100 DEG C ² Preferably at least 150J/cm ² More preferably at least 300J/cm ² . The steel has a minimum bending radius in the longitudinal or transverse direction of 5.0t or less, preferably 3.0t or less, more preferably 0.5t, and wherein t is the thickness of the steel strip or sheet.

Improved mechanical properties can be maintained even after post-weld heat treatment of the steel for 1 to 8 hours at a temperature in the range of 500 to 680 c, preferably for 4 to 8 hours at a temperature in the range of 600 to 640 c.

The following examples further describe and demonstrate embodiments within the scope of the present invention. These examples are given solely for the purpose of illustration and are not to be construed as limitations of the present invention, as many variations thereof are possible without departing from the scope of the invention.

Example 1

Table 1 lists the chemical compositions used to produce the panels tested.

Table 1 chemical composition (wt.%) of example 1.

	C	Si	Mn	Al	Nb	Cu	Cr	Ni	Ti	Mo	V
												Target object	0.035	0.4	1.55	0.03	0.03	0.25	0.2	0.15	0.015	0	0
Minimum value	0.025	0.3	1.48	0.02	0.025	0.15	0.1	0.1	0.005
												Maximum value	0.05	0.5	1.6	0.06	0.05	0.35	0.25	0.25	0.03	0.07	0.03

The test plate is prepared by a method comprising the following steps:

-heating to a temperature of 1140 ℃;

-hot rolling, wherein the controlled rolling compression ratio is 2.5, the final rolling temperature being in the range 840 ℃ to 880 ℃;

-accelerating the continuous cooling to about 100 ℃; and

-tempering at about 640 ℃.

Microstructure of microstructure

The microstructure can be characterized by SEM micrographs and the volume fraction can be determined using point counting or image analysis methods. The microstructure of the test plate includes 40% to 80% of quasi-polygonal ferrite, 20% to 40% of polygonal ferrite, and 20% or less of bainite, with the remainder being pearlite and martensite.

Yield strength of

Yield strength was determined according to ASTM E8 standard using a transverse specimen of a production lot of 2000 ton board. Average value of transverse yield strength (Rp _0.2 ) 508.+ -. 12MPa (FIG. 1).

Tensile strength of

Tensile strength was determined according to ASTM E8 standard using a transverse specimen of a production lot of 2000 ton board. The average value of the transverse ultimate tensile strength (Rm) was 590.+ -.1 MPa (FIG. 2).

Elongation percentage

Elongation was determined according to ASTM E8 standard using a transverse specimen of a production lot of 2000 ton board. Total elongation in transverse direction (A) ₅₀ ) The average value of (a) was 30.+ -. 1.4% (FIG. 3).

Flexibility of

The bending test involves plastic deformation of the test piece by using a single-stroke three-point bending until a specified bending angle of 90 ° is reached after unloading. The inspection and evaluation of the bending is a continuous process throughout the test series. This makes it possible to decide whether the punch radius (R) should be increased, maintained or decreased. If the same punch radius (R) is used in both the machine and transverse directions to meet a minimum bending length of 3m (without any defects), the material's bendability (R/t) limit can be determined in the test series. Cracks, surface necking marks, and flat bends (pronounced necking) are recorded as defects.

According to the bending test, the minimum bending radius (Ri) of the sheet in the longitudinal and transverse directions is 0.5 times the sheet thickness (t), i.e., ri=0.5 t.

PWHT resistance

Excellent tensile properties, such as a yield strength of at least 415MPa and an ultimate tensile strength of at least 550MPa, are maintained even after a severe PWHT treatment of 8 hours at 620 ℃.

Charpy-V impact toughness

Impact toughness values at-45℃were obtained by the Charpy-V notch test according to ASME (American society of mechanical Engineers) standards.

FIG. 4 shows that the average impact toughness value measured using a 6.7mm by 10mm transverse specimen of a 2000 ton plate from a production lot is 274J.

Fig. 5 shows the charpy-V impact toughness results for panels having different thickness in the machine and transverse directions. Table 1-1 summarizes the Charpy-V impact toughness results for panels of different thickness in the transverse direction.

TABLE 1-1 Charpy-V impact toughness of plates of different thickness

Thickness (mm)	KV(J/cm ² )	Temperature (. Degree. C.)	Direction
				10	338	-100	Transverse direction
20	587	-80	Transverse direction
				30	583	-60	Transverse direction
41	573	-60	Transverse direction

In the transverse direction, a test panel having a thickness of 10mm had an impact toughness of 338J/cm at a temperature of-100 DEG C ² The method comprises the steps of carrying out a first treatment on the surface of the A test panel having a thickness of 20mm had an impact toughness of 587J/cm at a temperature of-80 DEG C ² The method comprises the steps of carrying out a first treatment on the surface of the A test panel having a thickness of 30mm had an impact toughness of 583J/cm at a temperature of-60 DEG C ² The method comprises the steps of carrying out a first treatment on the surface of the A test panel having a thickness of 41mm had an impact toughness of 573J/cm at a temperature of-60 DEG C ² 。

Weldability of

The solderability test was performed on 41mm thick plates. The solderability test was performed by soldering three butt joints using a test piece of size 41mm x 200mm x 1000mm. Test pieces were cut from the plate in the main rolling direction so that 1000mm long butt welds were parallel to the rolling direction. The joint was welded using a 0.8kJ/mm heat input using flux-cored arc welding FCAW process No.136, and using a 3.5kJ/mm heat input using a single wire submerged arc welding process No. 121. The pre-heat temperature before the welding of the plates is in the range of 125 ℃ to 130 ℃ and the interlayer temperature is in the range of 125 ℃ to 200 ℃. The butt joint was welded using a half V groove preparation with a groove angle of 25 °. The selected welding material that may be consumed for the FCAW process is Esab Filarc PZ6138 with the EN/AWS classification T50-6-1Ni-P-M21-1-H5/E81T1-M21A8-Ni 1-H4. Selected welding materials that may be consumed for the SAW process are Esab OK Flux 10.62 and Esab OK Autrod 13.27 welding wire with the EN/AWS classification S-46-7-FB-S2Ni2/F7A10-ENi2-Ni 2. Welds welded by heat input of 3.5kJ/mm were tested under both weld conditions and PWHT conditions. The application of PWHT was performed at a temperature of 600℃for a holding time of 4 hours.

Tables 1-2 summarize the mechanical test results for the following welded joints:

-two lateral tensile tests of a rectangular specimen;

the charpy-V-shape test of the inclined plane was performed at-40 ℃ and-50 ℃ with three 10mm x 10mm samples from the following positions: fusion line +1mm (FL+1) and fusion line +5mm (FL+5); and

vickers hardness HV10 cross-weld hardness profile.

The mechanical test results show that the steel samples have excellent weldability at low temperatures and excellent HAZ toughness.

HIC resistance

HIC testing was performed according to NACE (American society of Corrosion Engineers) TM 0284. Fig. 6 shows the NACE TM 0284HIC test results for plates with different thicknesses. The test panels all showed an average (avg.) Crack Length Ratio (CLR) of less than 15%, indicating excellent performance of the steel in acid gas environments. The symbol "CSR" refers to the crack sensitivity ratio. The symbol "CTR" refers to the crack thickness ratio.

Example 2

Table 2 lists the chemical compositions used to produce the test panels. The slab number C002 is a comparative example.

The test panels were prepared by the method described in example 1.

The Final Rolling Temperature (FRT) and the cumulative compression ratio of the Controlled Rolling (CR) passes below the austenite non-recrystallization temperature are the main parameters that determine microstructure and mechanical properties. Table 2-1 summarizes the measured plate thickness, FRT and CR compression ratio. The slab numbers C002-1 and C002-2 are comparative examples.

Table 2-1 summary of measured plate thickness, FRT and CR compression ratios

Plate blank numbering	Thickness (mm)	FRT(℃)	CR compression ratio
				55106261	25	820	1.8
55106262	25	820	1.8
				55106331	25	800	3.0
55106031	41	820	1.8
				55106032	41	820	1.8
55106012	41	800	2.5
				55106049	41	850	3.0
E002-1	41	838	3.0
				C002-1	41	798	1.8
C002-2	41	777	2.5

Tensile properties

Tensile properties were determined according to ASTM E8 using transverse, 40mm wide and rectangular specimens. Fig. 7 shows that under transport conditions, all measured plates with a thickness of 10mm to 41mm have a yield strength higher than 480MPa and an ultimate tensile strength higher than 550MPa. The conveying conditions were defined as TMCP-ACC-T conditions without any further treatment after the steps of Accelerated Continuous Cooling (ACC) and tempering (T) in the thermo-mechanical control process (TMCP) used to manufacture the test panels of example 2. Post-weld heat treatment (PWHT) performed at 600 ℃ for 4 hours had little effect on tensile properties (fig. 7).

Full thickness tensile tests were performed on plates of thickness 12mm, 25mm or 41 mm. The greater the percentage of cross-section reduction before failure, the greater the ductility of the steel in the Z-direction. Fig. 8 shows that the percentage reduction in cross-sectional area is 77.6% to 81.8%, much greater than 35% required for standard grade ASTM a537CL 2.

Charpy-V impact toughness

Impact toughness was determined according to ASTM E23 using a 7.5mm 10mm longitudinal plate having a thickness of 10mm and a10 mm longitudinal plate having a thickness of 15mm, 20mm, 25mm or 41 mm. As shown in fig. 9, the charpy-V impact toughness is different for plates of different thickness. Table 2-2 summarizes the Charpy-V impact toughness results for panels of different thickness in the machine direction.

TABLE 2 Charpy-V impact toughness for different thickness plates

Thickness (mm)	KV(J/cm ² )	Temperature (. Degree. C.)	Direction
				10	>300	-68	Longitudinal direction
15	375	-68	Longitudinal direction
				20	300	-60	Longitudinal direction
25	375	-60	Longitudinal direction
				41	320	-52	Longitudinal direction

The impact toughness level of the 10mm thick and 15mm thick plates is at greater than 300J/cm at-68 DEG C ² In the upper shelf of (2), the energy of a 15mm thick plate under conveying conditions was 375J/cm ² . The impact toughness levels of the plates having a thickness of 20mm and 25mm under transport conditions or PWHT conditions at-60℃were 300J/cm, respectively ² And 375J/cm ² . The impact toughness level of 41mm thick plaques is 320J at-52 ℃.

FIGS. 10 and 11 show the effect of controlled rolling compression ratios on impact toughness of 25mm and 41mm thick plates (Table 2-1), respectively. Fig. 10 shows that increasing the controlled rolling compression ratio from 1.8 to 3 in a 25mm thick plate reduces the transition temperature from-52 ℃ to-60 ℃. In 41mm thick plates, increasing the controlled rolling compression ratio from 1.8 to 2.5 reduced the transition temperature from-40 ℃ to-60 ℃ (fig. 11). The best results were obtained when the controlled rolling compression ratio was 3.0 (fig. 10 and 11).

PWHT resistance

Post-weld heat treatment (PWHT) at 600 ℃ for 4 hours had little effect on tensile properties such as yield strength, ultimate tensile strength and elongation (fig. 7) or charpy-V impact toughness results (fig. 9-11).

Flexibility of

The method described in example 1 was used to measure the bendability. The minimum bending radius of a plate having a thickness of 41mm was 0.49 times the thickness of the plate in both the longitudinal and transverse directions (ri=0.49 t).

Microstructure of microstructure

The microstructure was characterized using the method described in example 1. As shown in fig. 12, the microstructure of the steel having a thickness of 41mm (table 2-1) includes quasi-polygonal ferrite, and bainite.

The level of Controlled Rolling (CR) compression ratio and the Final Rolling Temperature (FRT) have an effect on the grain size. By combining the controlled rolling compression ratio of 3.0 and the final rolling temperature of 838 deg., the desired microstructure of E002-1 as shown in fig. 9 (a) can be obtained. Higher controlled rolling compression ratios create more ferrite grain start sites, thereby reducing grain size. When the applied final rolling temperature is lower than 800 ℃, for example 798 ℃ in the case of C002-1 [ fig. 9 (b) ] or 777 ℃ in the case of C002-2 [ fig. 9 (C) ], the grain size is larger than when the applied final rolling temperature is higher than 800 ℃ [ fig. 9 (a) ].

Example 3

Table 3 lists the chemical compositions used to produce the panels tested. The slab number C003 is a comparative example.

The test panels were prepared by the method described in example 1.

Table 3-1 summarizes the cooling parameters of the panels tested. The accelerated continuous cooling stop temperature had little or no effect on mechanical properties (Table 3-2). However, the accelerated continuous cooling stop temperature is an important parameter for determining the low temperature toughness (tables 3 to 3).

Rolling tests with interrupted accelerated cooling were performed on 41mm thick plates, which indicated that accelerated continuous cooling to a temperature below 230 ℃ was very important for low temperature toughness. Upon interruption of accelerated cooling at a temperature in the range of 250℃to 290℃in Table 3-1, the Charpy-V impact toughness rapidly deteriorated at a temperature of-60℃in Table 3-3.

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Claims

1. A high strength steel comprising a composition consisting of, in weight percent (wt.%):

the balance being iron and unavoidable impurities, wherein the steel has a microstructure comprising a matrix consisting of, in volume percent (vol.%):

and the steel has the following mechanical properties:

a yield strength of at least 400MPa,

an ultimate tensile strength of at least 500MPa,

at a temperature of-50 ℃ to-100 ℃ of at least 34J/cm ² Charpy-V impact toughness.

2. The steel product as claimed in claim 1 wherein the steel product contains nonmetallic inclusions having an average inclusion diameter size in the range of 1 μm to 4 μm and wherein 95% of the inclusions have a diameter of less than 4 μm.

3. The steel product according to claim 1 or 2, wherein the steel product is a strip or a plate, the thickness of the strip or the plate being in the range of 6mm to 65 mm.

4. The steel product as claimed in claim 1 or 2 wherein the steel product has a yield strength of at least 415MPa.

5. The steel product as claimed in claim 1 or 2 wherein the steel product has an ultimate tensile strength in the range of 500MPa to 690 MPa.

6. The steel product as claimed in claim 1 or 2 wherein the steel product has a minimum bending radius in the longitudinal or transverse direction of 5.0t or less and wherein t is the thickness of the steel strip or sheet.

7. The steel product as claimed in claim 1 or 2 wherein the steel product has been subjected to a post-weld heat treatment at a temperature in the range 500 ℃ to 680 ℃ for 1 to 8 hours.

8. A method for manufacturing the high strength steel material according to any one of claims 1 to 7, comprising the steps of:

-heating a steel blank having the composition according to claim 1 to a temperature in the range of 950 ℃ to 1350 ℃;

i. The steel slab is subjected to a first plurality of rolling passes at a temperature above the austenite non-recrystallization temperature;

subjecting the steel blank from step (ii) to a second plurality of controlled rolling passes at a temperature below the austenite non-recrystallization temperature, wherein the controlled rolling passes have a reduction ratio of at least 1.5, and wherein the final rolling temperature ranges from 800 ℃ to 880 ℃;

-accelerating the continuous cooling to a temperature below 230 ℃ at a cooling rate of at least 5 ℃/s; and

optionally tempering at a temperature in the range 580 ℃ to 650 ℃ for 0.5 hours to 1 hour.

9. The method of claim 8, wherein the cumulative compression ratio of the hot rolling is in the range of 4.0 to 35.